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THE metal iron has physical and chemical properties which are somewhat different from those of steels, but a knowledge of the pure metal is a useful starting point in studying the behavior of steels. Solid iron exists in two different crystalline forms, called alpha (a) and gamma (?) iron, the former being the ordinary form at room temperature. When heated above the "A2 point," 768 C (1414 F), the metal becomes nonmagnetic, and at the "A, point," 910 C (1670 F), it undergoes a change in crystalline form to become gamma iron. In the range 768 to 910 C, it is sometimes called beta (ß) iron to distinguish it from the magnetic alpha, but the two are identical in crystal structure. When gamma iron is further heated to 1400 C (2552 E), it changes into delta (d) iron, which is identical with beta or non- magnetic alpha iron. Pure iron melts (in an inert atmosphere) at 1537 C (2800 F); the melting point is generally lowered by alloys or impurities. PROPERTIES OF IRON AT HIGH TEMPERATURES Many of the properties of iron undergo abrupt changes at the transition temperatures.1 For example, in heating through the A, point, a sudden shrinkage in volume occurs. Similarly, melting is attended by sharp physical changes such as an increase in volume. Table 16-1 contains a summary of the measured physical properties of pure iron at several temperatures. Usually the properties change smoothly within the range of stability of a given form, but this is not always true. For example, the specific heat reaches a sharp maximum at the A2 point although no trans- formation occurs. (See Fig 14-1.) Solvent Powers of Liquid Iron. Liquid iron is capable of dissolving most of the other metals and many nonmetallic elements as well. Of the metals that are liquid at the melting point of iron, only lead, silver, and bismuth refuse to dissolve to an ap-

Jan 1, 1951

By Earl F. Hanson

Fires in heavily timbered mines are disastrous, involving danger to both life and property. Some mines have been completely ruined or so heavily damaged that reopening them would not pay. Though few mines are entirely free from danger of fire, the risk is minimized by proper care and good design of plant. The Tintic Standard mine is at Dividend, Utah, in the famous Tintic mining district, about 60 miles south of Salt Lake City. The development of the mine was begun in 1907 by the late Mr. E. J. Raddatz and his associates. From the beginning, the work was hampered by high rock temperatures and the presence of high concentrations of rock gases such as carbon dioxide and nitrogen. Development work was carried on with great perseverance for nine years, and resulted in the discovery of ore in 1916. The total production of the Tintic Standard mine to Jan. 1, 1936, amounted to 1,664,388 tons of ore, having a gross value of $61,996,719. General Layout of the Mine The No. 2 shaft, which is the main working shaft of the Tintic Standard, is within the main orebody and is connected with the five other shafts. Their respective positions with respect to the No. 2 shaft are as follows: North Lily, 2770 ft. northwest with drift connections on the 700 and 900-ft. levels; No. 3 shaft, 670 ft. east connected through the 700-ft. level; Eureka Standard, 4550 ft. south, connected on the 900-ft. level; No. 1 shaft, 1470 ft. southwest, with connecting drifts on the 900 and 1250-ft. levels; Eureka Lilly, 2000 ft. southwest, with connecting drift and raise through the 1250-ft. level. The Tintic Standard orebody consists essentially of one continuous oreshoot extending from the 600-ft. level to the 1450-ft. level, with a cross-sectional area of approximately 70,000 sq. ft. at both top and bottom, reaching a maximum area of 215,000 sq. ft. between the 1100

Jan 1, 1937

By Russell Franks, W. O. Binder

In recent years a considerable tonnage of the 18 per cent chromium and 8 per cent nickel steels has been used in lightweight high-strength structures that must be resistant to deterioration under atmospheric conditions. These steels are austenitic in character and their strength can best be increased by cold-working. Although a marked increase in tensile strength is obtained, claims have been made that the cold-worked steels do not possess fully satisfactory elastic propertics. Under certain conditions, however, this has caused no great concern, because, as explained by Ragsdale,1,2 he cold-rolled steels have been widely applied for construction of lightweight, high-strength trains using spot welding for joining purposes. Ffield3 agrees that the ease with which thin sections of the cold-rolled 18-8 steels can be joined by spot welding is greatly in their favor, but that they would be more satisfactory if higher elastic properties and a higher modulus of elasticity could be obtained. He states that, as far as the mechanical properties are concerned, the straight chromium stainless steel is superior to the cold-worked 18-8 steel, and in order to illustrate, he prcsented the data of Fig. 1, which also give a stress-strain curve for a low-carbon 13 per cent Cr steel. Ffield† proposed that columbium-bearing or titanium-bearing cold-rolled 18-8 steels be given a heat-treatment (preferably for 1 hr.) in the range 400" to 600°C. to precipitate carbides in order to obtain better mechanical properties. Discussions by Sutton,4,5 de Ganahl,6,7,8 Krivobok and Lincoln,' Knerr,l0 and Thaden11 have given further information on the properties of the cold-rolled 18-8 steels compared with other materials used in lightweight, high-strength structures. These investigators and others point out that the ultimate tensile strength of the cold-rolled 18-8 steel can be increased to 200,000 lb. per sq. in. or higher; Krivobok et al. indicate that

Jan 1, 1940

By Russell Franks, W. O. Binder

In recent years a considerable tonnage of the 18 per cent chromium and 8 per cent nickel steels has been used in lightweight high-strength structures that must be resistant to deterioration under atmospheric conditions. These steels are austenitic in character and their strength can best be increased by cold-working. Although a marked increase in tensile strength is obtained, claims have been made that the cold-worked steels do not possess fully satisfactory elastic propertics. Under certain conditions, however, this has caused no great concern, because, as explained by Ragsdale,1,2 he cold-rolled steels have been widely applied for construction of lightweight, high-strength trains using spot welding for joining purposes. Ffield3 agrees that the ease with which thin sections of the cold-rolled 18-8 steels can be joined by spot welding is greatly in their favor, but that they would be more satisfactory if higher elastic properties and a higher modulus of elasticity could be obtained. He states that, as far as the mechanical properties are concerned, the straight chromium stainless steel is superior to the cold-worked 18-8 steel, and in order to illustrate, he prcsented the data of Fig. 1, which also give a stress-strain curve for a low-carbon 13 per cent Cr steel. Ffield† proposed that columbium-bearing or titanium-bearing cold-rolled 18-8 steels be given a heat-treatment (preferably for 1 hr.) in the range 400" to 600°C. to precipitate carbides in order to obtain better mechanical properties. Discussions by Sutton,4,5 de Ganahl,6,7,8 Krivobok and Lincoln,' Knerr,l0 and Thaden11 have given further information on the properties of the cold-rolled 18-8 steels compared with other materials used in lightweight, high-strength structures. These investigators and others point out that the ultimate tensile strength of the cold-rolled 18-8 steel can be increased to 200,000 lb. per sq. in. or higher; Krivobok et al. indicate that

Jan 1, 1940

By B. H. Kear, S. M. Copley

A dynamic theory of precipitation hardening in alloys containing cohevenl, stvess-.free, ordered particles is developed. The yield stvess is predicted from the stress dependence of the plastic strain vale, which is related to a derived dislocation velocity-stress function. The theory is shown to predict the magnilude and temperature dependence of the yield stress of a commercial nickel-base superalloy as well as many details of its slip structure. GLEITER and Hornbogen have developed a theory of the yield stress in alloys hardened by coherent, stress-free ordered ptrticles based on the observation that dislocations (b = a/2 (110)) glide in pairs.' When a dislocation pair shears through an ordered particle, the glide motion of the leading dislocation is impeded since it must create antiphase boundary (APB), whereas the motion of the trailing dislocation is assisted because it annihilates APB produced by the first one. Taking into account the APB energy, the size, and spacing of the particles, G-H predict the observed yielding behavior in Ni/18.8 Cr/6.2 A1 alloys composed of a matrix phase y '(nickel-solid solution) and particles of y' (basically Ni3A1). The G-H theory, as well as all previous theories of hardening by coherent ordered particles, do not include the details of dislocation motion; the present dynamic theory is an attempt to overcome this limitation. In the dynamic theory, the yield stress is predicted from the stress dependence of the plastic strain rate, which is related to a derived dislocation velocity-stress function. This function can be evaluated if the structure of the alloy, the APB energy of the ordered phase, and the characteristic drag or frictional stresses of the matrix and precipitate phases are known. The theory is shown to account for the observed temperature dependence of the yield stress in a commercial nickel-base superalloy. 1) NICKEL-BASE SUPERALLOYS Although their compositions are complex, the nickel-base superalloys have remarkably simple structures. A detailed examination of the commercial superalloy

Jan 1, 1968

By Josef Intrater

A dilatometric determination of the a + ß region temperature limits was performed on vacuum melted Zircaloy-2. The temperature of transformation for a-a+ß and a+ß — ß on heating and cooling as a function of rate of heating and heat treating history was studied. The limits were found to be 815° to 830°C to 975° to 995°C on heating and from 960° to 930°C to 785° to 770°C on cooling. A 5° to 15°C increase was observed for a- a + ß with increased rate of heating and a 20° to 30°C decrease for ß - a + ß with increased rate of cooling. 1 HE transformation temperature for pure zirconium has been determined by many investigators1-5 using metallography, electrical resistivity measurements, and dilatometry. The effect of the rate of cooling on the transformation temperature of zirconium has been studied by Hayes and Kaufmann6 and Duwez,7 and the effect of alloying elements by pfei18 and McGeary.13 Since the development of Zircaloy-2, the limits of the a + ß region for that alloy have been also determined. The first reported temperature for the a — a + a transformation in Zircaloy-2 was 900°C and for a + ß - ß 1000C Subsequent studies by Picklesimer and damson" indicated the beginning of the transformation at 810°C, and the metallographic study by the Technology Group of the Special Materials Section at the Westinghouse Bettis plant1' suggested that 820°to 830°C is the temperature of transformation. The dilatometric investigation by the latter group yielded 830°C for the beginning of the transformation. Thomas and Forscher12 set the beginning of the transformation at 900 °C on heating and the end of the transformation at 870°C on cooling. The above investigators stated in their paper that the dilatometric curve exhibits features caused by a eutectoid reac-

Jan 1, 1962

By Franklin H. Wilson

COPPER base alloys containing various amounts of aluminum and silicon are of considerable commercial interest. In particular the alloy containing 7 pct aluminum and 2 pct silicon shows an attractive combination of high strength with excellent machinability. Since there is no evidence in the literature that this ternary system has been studied since the recognition by Smith1 and Anderson2 in 1939 of the kappa phase of the binary copper-silicon system, the determination of the equilibrium phase relationships in the copper corner of the ternary diagram has been undertaken to provide a fundamental basis for understanding the behavior of these alloys. Phase boundaries have been determined by the micrographic examination of quenched samples while points on the liquidus surface have been obtained by thermal analysis. The results, of sufficient precision to serve their major purpose of providing a guide for the control of the manufacture of alloys in the region covered, are presented in a series of both isothermal and vertical sections of the diagram. ALLOYS STUDIED For a rough survey of the copper corner of this ternary system the compositions listed in Table I (Series A) were employed. The results of examinations of these alloys indicated that the compositions listed in Table 2 (Series B) would further clarify regions of greater complexity. All alloys were prepared as 1200 g heats, the copper being melted under charcoal, and then first silicon and then aluminum were added. An [ ] excess of 0.05 pct silicon was allowed to cover deoxidation losses. Two small bars 5/8 by 1 ¼ by 1 ¼ in. were cast using Durville type casting equipment. Analyses were made on all alloys using material which was hot-rolled at 800 or 850°C and annealed 17 hr at 650°C. All analyses were found to be within 0.05 pct of the nominal composition. Since the precision of analysis is not much better than this, the points of both isothermal sections and vertical sections at fixed aluminum and silicon contents have been plotted at the compositions indicated by the nominal composition.

Jan 1, 1948

By I. N. Knapp

I Herewith present some notes on the use of Portland cement to cement in the casing, and for plugging, to exclude water from oil and gas wells, and the methods employed. I have used my best efforts to. make each step of the operation of cementing wells perfectly clear. The information is the result of actual experience and observation. The illustrations are from original drawings made especially for this paper. Portland Cement It is not necessary to discuss here the manufacture of cement. Any established brand of the slower setting Portland cements may be purchased with the confidence that it will harden, if properly used, and exclude water from and plug wells efficiently. Setting.—By the setting of cement is meant its initial change from a soft or plastic mortar to a friable solid. After the cement has become thoroughly set it is still very weak and can be readily pulverized in the fingers. Setting is thought to be effected by the crystallizing out of the silicate and the aluminate of lime, which are soluble in water in their anhydrous form. After dissolving in the water they pass to the hydrated state, in which they are insoluble, and hence are precipitated in a crystalline form. Ordinarily neat Portland cement mortar remains perfectly plastic for from 1 to 1 1/2 hr. after mixing with water at about 70" F. A higher temperature accelerates and a lower retards setting.. Hardening.—This is due to a continual crystallization of salts from solution, and to further chemical and physical changes which develop slowly, but which continue for long periods of time. Testing.—It is no more necessary for an operator to go into the scientific testing of cement to be used in a well than to so test the pipe used. He must depend for the quality of each on the tests and reliability of the manufacturer. Cement or pipe may be damaged in shipment so

Jan 1, 1915

By William E. Ford, Edward Salisbury Dana

A. Hexagonal Division 1. Normal Class (13) Beryl Type 2. Hemimorphic Class (14) Zincite Type 3. Pyramidal Class (15) Apatite Type 4. Pyramidal-Hemimorphic Class (16) Nephelite Type 5. Trapezohedral Class (17) B. Trigonal or Rhombohedral Division 1. Trigonal Class (18) Benitoite Type 2. Rhombohedral Class (19) Calcite Type 3. Rhombohedral-Hemimorphic Class (20) Tourmaline Type 4. Trirhombohedral Class (21) Phenacite Type 5. Trapezohedral Class (22) Quartz Type 6. Other Classes (23) 7. Other Classes (24) Mathematical Relations of the Hexagonal System 119. The HEXAGONAL SYSTEM includes all the forms which are referred to four axes, three equal horizontal axes in a common plane intersecting at angles of 60") and a fourth, vertical axis, at right angles to them. Two sections are here included, each embracing a number of distinct classes related among themselves. They are called the Hexagonal Division and the Trigonal (or Rhombohedral) Division. The symmetry of the former, about the vertical axis, belongs to the hexagonal type, that of the latter to the trigonal type. Miller (1852) referred all the forms of the hexagonal system to three equal axes parallel to the faces of the fundamental rhombohedron, and hence intersecting at equal angles, not 90". This method (further explained in Art. 169) had the disadvantage of failing to bring out the relationship between the normal hexagonal and tetragonal types, both characterized by a principal axis of symmetry, which (on the system adopted in this book) is the vertical crystallo raphic axis. It further gave different symbols to faces which are crystallographally identical. It is more natural to employ the three rhombohedra1 axes for tri- gonal forms only, as done by Groth (1905), who includes these groups in a Trigonal Syslem; but this also has some disadvantages. The indices commonly used in describing hexagonal forms are known as the Miller-Bravais indices, since the were adopted by Bravais for use with the four axes from the scheme used by Miller in the other crystal systems.

Jan 1, 1922

By C. A. Heiland

A. INTRODUCTION (C. A. HEILAND) I. PRINCIPLES OF TEMPERATURE COMPENSATION IN MAGNETIC INSTRUMENTS The principle underlying the majority of magnetic intensity variom-eters is a comparison of the force to be measured with another force of known magnitude. The known force may be (a) of a magnetic nature, such as magnetic fields produced by magnets or coils, or (b) of a mechan-ical nature, such as (1) the elastic forces of wires, or (2) the force of gravity, such as employed in bifilar magnetometers and magnetic balances. In instruments of such type, usually a magnetic system is employed, free to move about either a horizontal or vertical axis, on some sort of a suspension. If the earth's magnetic field alone were present, such a system would adjust itself into the direction of the field or that of its components. Owing to the simultaneous action of the comparison force, however, the magnetic needle assumes a position that differs from the direction of the magnetic force; for all such instruments, the angle of deflection a is given by tan a = F/X, where F is the known and X the unknown earth magnetic force component. In other words, the magnetic system merely serves as an "indicator" for the position of balance of the two forces. Considering the matter from this point of view, it is not immediately obvious why the temperature of such a magnetic system should modify the results. In magnetic instruments, we may distinguish two different effects of temperature: (1) a magnetic effect and (2) a mechanical effect. The first is due to the fact that the magnetic moment of a magnet drops with an increase in temperature. This drop is characterized by the temperature coefficient of the magnetic moment µ (see equation 1). The second effect is due to the change in the .torsion of suspension wires with temperature or to the expansion of lever arms of masses acting on magnetic balances.

Jan 1, 1932

By Zay Jeffries

THIS paper will include a brief discussion of Prof. Howe's paper on The Supposed Reversal of Inheritance of Ferrite Grain Size from that of Austenite.1 The general subject of grain refining in steel and iron will also be treated. The conclusions reached are stated in the following hypotheses which are discussed more in detail below. 1. The ferrite grain size in pure iron, the ferrite and pearlite grain size in hypoeutectoid steel, the pearlite grain size in eutectoid steel and the cementite and pearlite grain size of hypereutectoid steel are not inherited from the grain size of the mother austenite. 2. The only structural feature that is generally inherited from the austenite of hypo- and hypereutectoid steels on cooling through their transformation ranges is the position of the excess ferrite or cementite at the austenite grain boundaries, sometimes causing complete and sometimes incomplete networks which outline the old austenite grain boundaries. Rapid cooling through the transformation range will cause the non-inheritance of this structural (network) feature. 3. The austenite grain boundaries themselves are nearly always effaced in all steels and also in pure iron during the Ar transformations. 4. The grain-size refining of steel and iron is brought about by the combined. effect of non-inheritance of the transformation products on either heating or cooling, i.e., the austenite transformation .products do not inherit their grain size from the austenite on cooling through the transformation range for does austenite inherit its grain size from the products which form austenite on heating. 5. In general, in both iron and carbon steel the larger the austenite ,grain size, the larger will be the grain size of the transformation products on cooling. This, of course, assumes all other conditions constant except the austenite grain size. An exception is found to this general rule in very pure iron such as electrolytic iron. 1n this instance small austenite grains may form very large ferrite grains on cooling through Ar3.

Jan 11, 1917

By L. B. Slichter

It has been estimated that rock exposures in most mining districts aggregate less than 1 per cent. of the total surface area.1 Conclusions concerning the hidden 99 per cent. necessarily have been based on knowledge of only a small fraction of the surface, and from this limited information the geologist has been obliged to deduce the approximate and general nature of the underlying formations. In the search for ore it is necessary to supplement this general picture with greater detail. Fairly accurate knowledge of relatively small zones is required. This knowledge is carefully sought in many different ways—detailed interpretations of geological history, the tracing of float, the use of test pits, of trenching, of drilling, and of other methods. Practical ingenuity, scientific brilliance, great perseverance and patience, and the expenditure of large sums of money have all been a part of this search. The problem of exploring beneath cover with economy, and of obtaining essential information which this cover hides, is obviously a difficult one. It is also a very important one, for geologist's inform us that already the story of the 1 per cent. exposed outcrops has been largely read.V here remains the nearly virgin field comprised in the hidden 99 per cent.: and the importance of this zone is bound to increase continually. In recent years, the problems connected with the exploration of this "99 per cent. zone" have attracted a constantly enlarging group of investigators and at present greater efforts than ever before are being made to resolve them. This work, in it's various special aspects, has enlisted, beside geologists, the cooperative efforts of chemists, physicists, and mathematicians. The progress made in this development should be kept before those directing explorations in order that due advantage may be taken of a changing situation. My subject concerns certain aspects of one of the older physical methods of exploration, the magnetic method. This method is familiar in principle to most of us, arid because of its age., is especially suitable for

Jan 1, 1929

By John P. Harmon

This paper was written with two objects in view: (1) To describe in detail what has been done toward the ventilation of the main unit of the Portovelo mines and the results; (2) to give information that may be useful to anyone who may have to make a preliminary mine ventilation survey, with recommendations for the improvement of the existing system and finally the maintenance of the improved system. General Description The mining operations of the South American Development Co. at Portovelo, Province of El oro, Ecuador, are described here only briefly, as a full discussion does not fall within the scope of this paper. The orebodies are of the steeply dipping, gold quartz vein type, averaging about 2.0 meters wide. As shown in Fig. 1, the ore chutes are relatively far apart, making considerable development work necessary before connecting raises can be driven between the levels. Faulting' on a large scale and a parallel vein system have made long dead-end drifts and many crosscuts necessary in the exploration of veins. The ventilation of these dead ends is described under Auxiliary Ventilation. The method of mining2 is the filled rill stope method, similar to the method that is largely employed at the Butte mines. The main shaft of the mine is in the Amarillo Valley, on either side of which mountains rise abruptly to a height of several hundred feet above the collar of the shaft. Entrance is made to the mine: (1) Through the American shaft, which extends vertically from the surface to a depth of 1100 ft., of which only the first 700 ft. are in active use at the present time; (2) levels are driven from the American shaft at intervals of 30 m. (100 ft.), of which only the third, fifth and seventh levels need be considered in the ventilation problem; (3) adits at intervals of 30 m. (100 ft.), with a few exceptions, above the collar of the American shaft.

Jan 1, 1931

By W. R. Hibbard

The term lineage was first introduced by Buergerl to denote dendritic branches, grown from a crystal nucleus during solidification from the liquid, with imper- fections in alignment of the order of 10 : to 10-4 cm. These imperfections are associated with distortions of the dendrites caused by the subsequent solidification of the inter-dendritic filling and are illustrated in Fig I. Buerger etched zinc single crystals made by the Bridgman3 method to develop this dendritic pattern and show lineage microscopically. Since maximum crystal growth in face-centered cubic metals occurs along the [IT] directions,2 lineage would appear in copper as small deviations of dendritic branches from these directions. Elam4 and others5"8 reported etching effects resulting from coring along the cube face traces of various copper alloy crystals. These traces were used by Samans8 to orient single crystals. Williamson9 noted a "lined structure (in alpha brass single crystals) consisting of an ordered arrangement of etch pits" which were probably traces of {IOO} planes but somewhat uncertain because they were never straight (that is, apparently lineage). Greningerl1 described the distortion and multiplicity of Laue spots in back reflection X ray photograms of copper single crystals showing lineage. The present investigation -includes metal-lographic, X ray, and property studies of lineage in large grain and single crystal specimens of copper. Preparation of Specimens specilnens 10 in. long 35 1/2 in. diam of oxygen-free copper (o.F.H.c.) (99.985 pet) were' melted vertically in graphite crucibles under an atmosphere of commercial nitrogen purified through a gas train including sulfuric acid, potassium chloride, magnesium chiorate, hot copper chips and phosphorus pentoxide. After a temperature of 1200°C had been maintained for 1/2 hr, the furnace was turned off and the specimens cooled at a rate of about 150°c an hour' This method of direct solidification was used by Greningerlo to produce single crystals with marked lineage and in this investigation produced a fair yield of single crystal specimens in

Jan 1, 1949

By W. R. Hibbard

The term lineage was first introduced by Buergerl to denote dendritic branches, grown from a crystal nucleus during solidification from the liquid, with imper- fections in alignment of the order of 10 : to 10-4 cm. These imperfections are associated with distortions of the dendrites caused by the subsequent solidification of the inter-dendritic filling and are illustrated in Fig I. Buerger etched zinc single crystals made by the Bridgman3 method to develop this dendritic pattern and show lineage microscopically. Since maximum crystal growth in face-centered cubic metals occurs along the [IT] directions,2 lineage would appear in copper as small deviations of dendritic branches from these directions. Elam4 and others5"8 reported etching effects resulting from coring along the cube face traces of various copper alloy crystals. These traces were used by Samans8 to orient single crystals. Williamson9 noted a "lined structure (in alpha brass single crystals) consisting of an ordered arrangement of etch pits" which were probably traces of {IOO} planes but somewhat uncertain because they were never straight (that is, apparently lineage). Greningerl1 described the distortion and multiplicity of Laue spots in back reflection X ray photograms of copper single crystals showing lineage. The present investigation -includes metal-lographic, X ray, and property studies of lineage in large grain and single crystal specimens of copper. Preparation of Specimens specilnens 10 in. long 35 1/2 in. diam of oxygen-free copper (o.F.H.c.) (99.985 pet) were' melted vertically in graphite crucibles under an atmosphere of commercial nitrogen purified through a gas train including sulfuric acid, potassium chloride, magnesium chiorate, hot copper chips and phosphorus pentoxide. After a temperature of 1200°C had been maintained for 1/2 hr, the furnace was turned off and the specimens cooled at a rate of about 150°c an hour' This method of direct solidification was used by Greningerlo to produce single crystals with marked lineage and in this investigation produced a fair yield of single crystal specimens in

Jan 1, 1949

By L. W. Eastwood

The structure of eutectics has been studied by a number of investigators, and the complexity of the structural relationship of the components has been agreed upon, especially that of the "eutectic colony," type III of Portevin's1 classification. This paper does not attempt to alter the present classifications of eutectics; rather, it discusses the relationship between the macrostructure and the microstructure of the copper-cuprous oxide eutectic, and describes in detail the mechanism of the formation and the structure of the eutectic. Further, throughout the entire discussion it attempts to point out certain misconceptions in regard to the interpretation of the observed structure of the copper-cuprous oxide eutectic, one of the so-called "eutectic colony" types, and to correlate as much as possible the work done by the writer with that done by other investigators. Of the three types of eutectics listed by Portevin,2 the eutectic colony type has presented the greatest difficulty in interpretation. Although the copper-cuprous oxide eutectic is typical of the eutectic colony type in appearance, it may not be truly representative because the copper forms about 96 per cent of the eutectic and the cuprous oxide about 4 per cent. The interpretation of the origin and structure of the "eutectic colony" given here is strictly limited, therefore, to the copper-cuprous oxide eutectic. Macrostructure Fig. 1 shows the macrostructure of the eutectic at 3 diameters, the melt having been chill-cast into an iron mold, the resulting specimen sawed into two parts and the section polished and etched with the usual mixture of hydrogen peroxide and ammonium hydroxide. This specimen contains 0.48 per cent oxygen and is only slightly hypereutectic. If such a melt is allowed to cool slowly, large macrograins will form without the columnar structure of Fig. 1 and the well-known phenomenon of liquation will occur, forming copper oxide dendrites at the top and copper

Jan 1, 1934

By A. S. Eve

In the spring of 1930, the question was raised as to the possibility of estimating the depth to which the pyrrhotite-nickel deposit at the Falconbridge mine extended in the earth. This body is 7500 ft. long, it dips at 80° to the north, and it is fairly straight with slight sinuosities. Dr. F. W. Lee suggested that it might be possible to solve the problem by measuring the vertical magnetic intensity, with a good magnetometer, both near the surface of the earth and also at a level a few feet higher in the air. The writer regarded this suggestion with some scepticism, but it was decided to put the scheme to a practical test. In July, 1930, Dr. Lee, Dr. Keys and the writer, with Mr. Joyce and Mr. Belloc as assistants, made a firm, light wooden platform (Fig. 1) measuring 4 by 4 ft. on the top, which was 10 ft. above ground. This platform could be moved easily a few yards by four men, and carried across country in a wagon or by motor truck. Some iron nails were used in its construction, but we were not able to detect their influence on the Askania magnetometer. Nevertheless, in future work, the use of iron should be completely avoided. Readings were taken along a traverse at right angles to the orebody, at several stations, both on the ground and on the platform. The excess readings, over the dike, as compared with those about 800 ft. away, are shown in Table 1. The largest reading, 6.8, is open to suspicion and should be rejected. The depth of the top of the orebody is known by electrical survey, and by actual drilling, to be about 112 to 123 ft., and the covering is sandy glacial drift. Similar readings were taken over a long straight magnetic diabase dike. These also are shown in Table 1. The magnetic vertical intensity is consistently less at the higher level, which is 10 ft. above the lower level.

Jan 1, 1932

By Albert Sauveur

(Cleveland Meeting, October, 1912.) ALTHOUGH many metallurgists know that some pearlitic special steels can be made troostitic, martensitic, and even austenitic, without quenching, and, therefore, without exposing them to the dangers of the quenching-bath, the practical significance of such possibility does not seem to have received the attention it deserves. Let us recall briefly the mechanism of the transformation alluded to. In Fig. 1 (after Guillet) is shown the constitutional diagram of nickel-steel that; is, of alloys of iron, carbon, and nickel. Any steel the composition of which falls within the area LPO is pearlitic after slow cooling; any steel the composition of which is represented by a point within the area MLO is martensitic, and any steel within the area MNO is austenitic. Starting with pearlitic steel, then by increasing the carbon-content or the nickel-content, or both, the metal becomes martensitic and eventually austenitic; that is, its composition crosses first the boundary-line LO and then the line M0-and it remains martensitic or austenitic after slow (air) cooling.1 We have inn this way imparted to the metal without quenching it the properties of hardened carbon-steel. In other words, the steel is now self-hardening. Nickel cannot be introduced into solid steel, but by the well-known process of case-hardening it * Professor of Metallurgy and Metalllography in Harvard University. + Assistant in Metallurgy and Metallography in Harvard University. 1 The boundaries between the various zones are not, as a matter of fact, as sharp as indicated in the diagram, because pearlite is not transformed abruptly and in its entirety into martensite when a certain critical composition is attained, nor is martensite suddenly and completely converted into austenite. Greater refinement in the construction of the diagram would undoubtedly result in the introduction of troosto-pearlitic, troostitic, and troosto-martensitic zones between the pearlitic and martensitic areas and-of a martenso-austenitic zone between the martensitic and austenitic areas. Fur the purpose of this note, however, the existence of these transition-conditions may be ignored.

Nov 1, 1912

By Thomas C. Wilder

The concentration and activity of oxygen in liquid copper has been measured by the reversible galvanic cell: at 1100°and 1200°C for the concentration range of oxygen of 3 to 5600 ppm Measurements for approximately thirly alloys at each temperature indicate that the activity of oxygen roughly follows Henry's law, the activity coefficients of oxygen being 0.12 and 0.205 at 1100° and 1200°C, respectively. These results compare javorably with the results of seven other electrochemical and gaseous equilibrium studies on this system, the free energy of the solution reaction 1/2 O2 = O (7 at. pct in Cu, liq) being -18.10 * 0.30 kcal per gfw at 1200°C. Estimates of the standard molar free energy of formation of CupO(S) are also made. Praclicul consideration of this cell as an oxygen sensor for liquid metals is discussed briefly. SOLID mixed-oxide electrolytes which have a defect structure in the anion sublattice have recently become a very popular means by which the chemical potential of oxygen is measured at high temperatures. Since the pioneering work of Kiukkola and Wagner1 such electrolytes have been used in numerous cell studies for determining the thermodynamic properties of metallic oxides and oxy compounds, the partial pressure of oxygen in vacuum systems and other atmospheres, and the properties of oxygen dissolved in metals and alloys. The use of a solid electrolyte in a cell for a fast reliable method of oxygen determination in liquid metals would be a very important tool to the process metallurgist. In copper, for example, the oxygen content of the molten metal in the cathode furnace is a prime factor in determining the "set" or solidification profile of cast wire bars. In this investigation the activity of oxygen in liquid copper was measured by an electrochemical cell employing the electrolyte ZrO2 + 7 1/2 wt pct CaO. The cell used at 1100° and 1200°C was

Jan 1, 1967

By J. B. Berkowitz-Mattuck, R. Hopper, L. Kaufman, E. V. Clougherty

MARNOCH has reported on the attractive high-temperature oxidation resistance of Hf-Ta alloys.' In view of current interest in these alloys, three Hf-Ta alloys were tested in the present study. These included a Hf-27Ta alloy and a Hf-20Ta-2Mo alloy which were supplied by Marquardt Corp. and a Hf-20Ta-Zr alloy provided by Illinois Institute of Technology Research Institute.' The microstructure of these alloys was heated inductively in a glass vacuum system; the oxidizing gas, helium carrier plus a known partial pressure of oxygen (35-102 Torr), was continually passed over the sample at temperature for a given time period. The concentration of oxygen in the gas stream was continuously measured by a differential thermal conductivity technique2 so that the rate of oxygen picked up by the sample could be obtained. The flow rates were 0.1 to 0.2 ft per sec in the vicinity of the specimen. The concentration of oxygen did not fall below 50 pct of the input level in any of the experiments. In several cases a single sample was exposed consecutively at a series of increasing temperature levels. In a second system at ManLabs, cylindrical specimens were exposed to flowing air in a z irconia tube which was heated by radiation from a carbon tube resistance furnace. The specimen was heated from room temperature to the test temperature in flowing argon, then oxidized in air flowing at 0.1 to 0.2 ft per sec. Following exposure, the air was displaced by argon and the specimen was cooled. Postoxidation metallography was performed on every sample. The depth of conversion of the metal matrix to oxide was measured as shown in Fig. 1. The conversion depth was obtained by measuring the oxygen-contaminated metallic zone (which is bounded by a dense outer oxide and a subscale oxide region), subtracting its dimensions from the original dimensions, and dividing by two. Preferential oxidation of

Jan 1, 1968